Nano-precipitation strengthened cold-rolled batch annealed high strength low alloy steel sheet

ABSTRACT

A high strength low alloy steel sheet product produced by a continuous strip process. The sheet product comprises from 0.045 to 0.06 weight percent Carbon, from 0.75 to 1.2 weight percent Manganese, from 0.02 to 0.04 weight percent Aluminum, and from 0.075 to 0.12% weight percent Titanium and an Fe balance. The steel sheet product has been subjected to cold rolling and batch annealing to form a steel sheet product having a yield strength of at least 500 MPa, a substantially ferritic microstructure with nano TiC precipitates and a hole expansion ratio of at least 60%.

CROSS-REFERENCE TO RELATED APPLICATION

This application claims the benefit of U.S. Provisional PatentApplication No. 63/341,811 filed May 13, 2022, which is incorporatedherein by reference.

BACKGROUND

The present disclosure generally relates to a high-strength low-alloysteel having a yield strength of at least 490 MPa. The disclosureincludes a process for manufacturing a low carbon Ti-microalloyed ofsteel for the development of cold-rolled and batch-annealedhigh-strength high-formable steels. In some embodiments, the processincludes a clean steel practice combined with hot rolling and batchannealing utilized to retain precipitate strengthening in the finalfully processed steel and microstructural attributes that results inincreased forming characteristics.

Development of high strength low alloy steels (HSLA) has redefined steelusage in various industries, for example, construction, machinery,automotive, agriculture and transportation. The HSLA steel providesengineering as well as economic benefits such as, weight reduction,improved and excellent weldability, reduced cost of engineeringconstructions, safety of components, increased payloads fortransportation and enhanced fuel economy for passenger cars. Theautomotive industry has benefitted through advent of higher strengthHSLA steels as this steel has provided for increased vehicular safety,enhanced fuel economy, and minimized CO2 emissions. Automotivemanufacturers are now leaning towards use of greener HSLA steels in aneffort to minimize CO2 emission. Companies are requesting that steelmanufacturers provide carbon-neutral or carbon-minimal HSLA steels andthe global steel industries are gearing towards innovating steelprocessing ideas and technologies that provide sustainable solutions togreen steel production. In this evolving perspective, mini steel millsemploying continuous strip production (CSP) technology have significantadvantage in reducing CO2 emissions over discrete integrated steel millsthat runs on blast furnace ironmaking technology. Moreover, CSP millsintegrated with advanced steel processing technologies are competingwith thick slab casting discrete steel mills not only in productionthroughput but also in offering advanced high strength grades of steelswith lean alloying and least cost of production.

HSLA sheet steels with a yield strength of at least 490 MPa areincreasingly demanded for many stamped parts as well as tubular forms ofautobody applications. High yield strength HSLA steels have not beenwidely developed. In the past, the addition of vanadium (V) and/orniobium (Nb) microalloyed HSLA steels was investigated. A solid solutionwas strengthened as well as microalloy V, Nb strengthened sheet steelsand mostly with production through thick slab casters and with heavyalloying. Furthermore, these developments also received less attentionbecause of the emergence of dual-phase (DP) steels offering high tensilestrengths and strain hardening capacity. However, DP steels are heavilyalloyed and the ferrite yield strength is lower due to pre-yielding whencompared with similar-tensile strength HSLA steels.

SUMMARY

The present disclosure includes a high strength low alloy steel sheetproduct produced by a continuous strip process. The sheet productcomprises from 0.045 to 0.06 weight percent Carbon, from 0.75 to 1.2weight percent Manganese, from 0.02 to 0.04 weight percent Aluminum, andfrom 0.075 to 0.12% weight percent Titanium and an Fe balance. Whereinthe steel sheet product has been subjected to cold rolling and batchannealing to form a steel sheet product having a yield strength of atleast 500 MPa, a substantially ferritic microstructure with nano TiCprecipitates and a hole expansion ratio of at least 70%.

The present disclosure includes a method for producing a high strengthlow alloy steel sheet product comprising from 0.045 to 0.06 weightpercent Carbon, from 0.75 to 1.2 weight percent Manganese, from 0.02 to0.04 weight percent Aluminum, and from 0.075 to 0.12% weight percentTitanium and an Fe balance, a yield strength of at least 500 MPa, asubstantially ferritic microstructure with nano TiC precipitates and ahole expansion ratio of at least 60%, comprising the steps of:continuously casting a steel slab approximately 55 mm-85 mm thick;maintaining a temperature of the steel slab; hot rolling the steel slabto a steel sheet at a finishing temperature of 875-950C; cooling thesteel slab; coiling the steel sheet at a temperature of 600-675C; coldrolling the steel sheet to 40-75%; and hydrogen batch annealing thesteel sheet at temperature of 600-650C to achieve a fully recrystallizedferritic microstructure with nano precipitates of TiC.

The present disclosure includes a high strength low alloy steel sheetproduct produced by a continuous strip process. The sheet productcomprising from 0.045 to 0.06 weight percent Carbon, from 0.75 to 1.2weight percent Manganese, from 0.02 to 0.04 weight percent Aluminum,from 0.075 to 0.12% weight percent Titanium, 0.04 weight percent Niobiumand an Fe balance. Wherein the steel sheet product has been subjected tocold rolling and batch annealing to form a steel sheet product having ayield strength of at least 500 MPa, a substantially ferriticmicrostructure with nano TiC precipitates and a hole expansion ratio ofat least 60%.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 is a prior art graph showing the effect of Mn and carbon contenton the r m values of cold rolled annealed steels.

FIG. 2 is a prior art graph showing the relationship between the stateof austenite recrystallization during hot rolling and the transformedferrite grain size.

FIG. 3 is a prior art graph showing the influence of microalloying onthe recrystallization stop temperature in low carbon HSLA steels.

FIG. 4 is a graph showing the yield and tensile properties of coldworked samples of selected steels after simulated batch annealing.

FIG. 5 is a graph showing the variation in the strength properties ofTi- and Ti—Nb steels across the coil width after batch annealing.

FIGS. 6 a and 6 b are micrographs showing the optical microstructures ofTi steel and Ti—Nb steel after hot rolling.

FIGS. 7 a through 7 d show the optical and scanning electron micrographsof batch annealed, Ti-steel, and Ti—Nb steel showing fullyrecrystallized microstructures.

FIGS. 8 a and 8 b show inclusions found in the final fully annealedsheet steels revealing mainly globular oxide and oxy-sulfide inclusions.

FIGS. 9 a through 9 c are TEM images of hot rolled Ti—Nb steel showingcoarse precipitates along grain boundaries and an overview of fineprecipitates and some localized areas of precipitates in rowarrangement.

FIGS. 10 a and 10 b show the chemical composition of precipitates in hotrolled Ti—Nb steel.

FIGS. 11 a and 11 b shows the fine nano-sized precipitates withinferrite grains and some along the grain boundaries.

FIGS. 12 a through 12 c are images showing the precipitates in hotrolled Ti-Steel as examined in thin foils, distribution of fineprecipitates.

FIG. 13 shows HAADF image and EDS mapping of precipitates larger than 20nm were Ti, C and S bearing.

FIG. 14 shows an HAADF image and EDS mapping of nano-sized precipitatesin hot rolled Ti-steel indicating Ti as major element.

FIG. 15 shows a graph of EELS spectrum showing the presence of Ti andcarbon only in a fine precipitate.

FIGS. 16 a and 16 b show images of fine precipitates in ferrite inTi-steel and HAADF image and EDS mapping of precipitates in Ti-steelrevealing presence of Ti only in fine precipitates less than 10 nm.

FIGS. 17 a and 17 b show TEM images of fully recrystallizedmicrostructures of Ti and Ti—Nb steels.

FIGS. 18 a and 18 b show images of the microstructures of Ti-steel aftersimulated batch annealing at 630° C. and 700° C. revealing significantferrite grain growth after recrystallization.

FIGS. 19 a through 19 c show a grain boundary mapping of Ti-steel andTi—Nb steel constructed from orientation imaging from electron backscattered diffraction (EBSD).

FIGS. 20 a and 20 b show an ODF of batch annealed Ti-steel and Ti—Nbsteel showing ODF of batch annealed Ti-steel and Ti—Nb steel showingincreasing γ-fiber orientations.

FIGS. 21 a through 21 d are charts showing the precipitate sizedistribution in hot rolled and annealed Ti-steel.

DETAILED DESCRIPTION

Mechanical property requirements for HSLA 490 sheet steels are listed inTable 1 as required by various common sheet steel OEM specifications.Table 1 shows the tensile property requirements of HSLA 490 annealedsheet steels as required by various specifications.

TABLE 1 Tensile % Yield Strength Elonga- Strength, MPa MPa tionSpecification Grade min max Min min SAE J2340 490 X 490 590 560 14 SAEJ2340 490 Y 490 590 590 12 ASTM A 1008 HSLAS Grade 70 480 585 14 Class 1ASTM A 1008 HSLAS Grade 70 480 550 14 Class 2 ASTM A 1008 HSLAS-F Grade70 480 550 16

In some cases, a Ti-only alloy approach was utilized because very finesized TiC can be precipitated in hot rolled coils giving a significantincrease in ferrite yield strength. Ti is effective in scavengingnitrogen from steel and tying up with most of the carbon to form nanosized precipitates provided hot rolling parameters are properlycontrolled. The extent of precipitation strengthening can be significantand is approximated by a σ_(ppt)=B. (wt. % alloy). In some embodiments,B is an average 1500 MPa/wt. % Ti. Carbon was maintained below 0.06 wt.% as excessive carbon to that required for TiC formation will causeFe-carbide precipitation during coiling and may deteriorate drawabilityand stretchability. Solid solution elements such as Mn, Si aresignificant ferrite strengtheners as outlined in Eq [1], as set forthbelow. In some embodiments, Mn additions were kept low as an excess ofMn introduces mobile dislocations in ferrite during rapidtransformations after hot rolling through its influence on Ar₃temperature and such mobile dislocations interfere with yield strengththrough pre-yielding. Mn also adversely influences r_(m) value inannealed sheet steel as shown in FIG. 1 .

Total strengthening will thus be additive of ferrite grain size andprecipitation strengthening as given below in Eq. [1] as deduced from atreatise of strengthening mechanisms in ferrite.

σ_(y)(MPα)=88+37mn'83Si+2900N _(free)+17d ^(−1/2)+σ_(ppt)+σ_(d)  Eq. [1]

In some embodiments, d is grain diameter in mm, a σ_(ppt) is theprecipitation strengthening, σ_(d) is dislocation strengthening.

FIG. 1 shows the effect of Mn and carbon content on the values of coldrolled annealed steels. Grain size and strength of annealed sheet steelshave a direct inheritance from hot rolled steel as indicated in FIG. 2 .FIG. 2 shows the relationship between state of austeniterecrystallization during hot rolling and the transformed ferrite grainsize. A finer fully recrystallized austenite grain results in fineferrite grain after transformation. Since TiC has limited solubility inaustenite finish rolling is carried out at relatively higher temperaturethan Nb-microalloyed steels. Austenite is refined through repeatedrecrystallization rolling above non-recrystallization temperature (TNR)which is lower than that of Nb-bearing steels as shown in FIG. 3 fromsimilar steel. As seen from experimental data in FIG. 3 , Ti influenceon TNR diminishes in excess of 0.10 wt. % and the TNR can be estimatedto be around 917° C. An accelerated cooling after finish rolling alsoinduces many nucleation sites for precipitation as well as inducesnon-polygonal fine ferrite formation. Additionally, to examine theinfluence of small addition of Nb on the starting austenite grain sizeand therefore, the final ferrite grain size a second alloy with Nbaddition was considered. FIG. 3 shows the effects of microalloying onthe recrystallization stop temperature in low carbon HSLA steels (C≤0.07wt. %).

The present disclosure describes a low carbon, lean Ti-bearing alloy todevelop ferritic HSLA 490 sheet steel in cold rolled-batch annealedcondition. Strengthening in final sheet steel was obtained primarilythrough ferrite grain size and TiC precipitation strengthening withmoderate solid solution strengthening. Precipitation of TiCnano-precipitates in hot rolled steel was induced through control of thefinishing deformation temperature, cooling rate and coiling temperature.A batch annealing cycle appropriate of the current steel was evolvedthrough pre-simulations to determine a full recrystallization annealingtemperature. Actual mill production of sheet steels of variousthicknesses revealed HSLA 490 steels could be successfully developedwith excellent elongation and strain hardening index using a leanTi-only HSLA approach. Nb addition of up to 0.04wt. % to the chosenalloy resulted in similar properties but didn't enhance tensile orformability properties over those of Ti-steel. TEM studies ofmicrostructure and precipitation revealed fully recrystallized very fineferritic microstructure with homogenously distributed nano TiCprecipitates of 1.5-8 nm in size with an average 3.16 nm. Precipitatesize or distribution were not affected by the recrystallization annealtemperature. Hole expansion ratio values of more than 90% were obtainedfor 1.2 mm thick Ti-HSLA sheet steels. Nano-sized TiC precipitates, finerecrystallized grain size and excellent internal cleanliness contributedto the hole expansion ratio values. Formability as indicated throughstretchability indicator, hole expansion ratio far exceeds that reportedby similar strength sheet steels.

HSLA steels with yield strength of at least 490 MPa having excellentstretch forming capability will be described herein. In someembodiments, a predominantly ferritic Ti-microalloyed HSLA steel mayinclude a higher yield strength as well as formability throughnano-scale precipitation in ferrite. The nano scale precipitates whenprecipitated within ferrite not only increases strain hardening offerrite but also resists recrystallization during annealing.

In some embodiments, a very low carbon C—Mn based ferritic steel withTi-only microalloying was selected to develop a family of HSLA sheetsteels with a minimum yield strength of 490 MPa meeting SAE J2340 CR490X specifications. Batch annealing was utilized in place of aconventional continuous annealing approach. Batch annealing providesresults in higher total elongation and higher drawability compared tocontinuous annealing, batch annealing can be performed at lowertemperatures, i.e. just high enough to complete recrystallization ofcold deformed structure without disturbing the precipitate state, lowtemperature annealing prevents precipitation coarsening and grain growthis avoided, batch annealing also retains excellent flatness of thingauge sheet steels as compared to continuous annealing because ofprevention of heat buckling. Batch annealing with optimized cycle mayresult in better cross width property uniformity compared withcontinuous annealing.

Strength development is based primarily on precipitation strengtheningand control of ferrite grain size. Details of alloy design, hot rollingprocessing approach and annealing parameter optimization through studiesof microstructure and precipitation evolution at each stage ofprocessing are described herein.

Steelmaking and Processing

Chemistry of the HSLA 490 steel is shown in Table 2. A clone chemistrywith small Nb addition (<0.04 wt. %) was also chosen to examinepartially recrystallized austenite grains on the final ferrite grainsize and properties. Carbon was restricted to less than 0.06 wt. % (i)to minimize grain boundary iron carbide formation and allow onlysufficient amounts for micro-alloy carbide formation, (ii) to controlTi:C stoichiometry for nano scale precipitation of TiC precipitates.Table 2 shows the chemistry of steels selected in the current study (wt.% max).

TABLE 2 Steel C Mn P S Si Ti N Nb CE_(IIW) P_(cm) Ti 0.060 1.1 0.0150.003 0.30 0.12 0.009 Res 0.25 0.13 Ti-Nb 0.060 1.1 0.015 0.003 0.300.12 0.009 0.04 0.25 0.13 CE_(IIW): C + Mn/6 + (Cr + Mo + V)/5 + (Cu +Ni)/15; P_(cm): C + Si/30 + (Mn + Cr + Cu)/20 + Ni/60 + Mo/15 + V/10 +5B.

The steel was manufactured utilized an electric arc furnace ensuringleast slag carryover and low residuals through suitable choice of scrapmix. Deoxidation practice was aimed at low total dissolved oxygen sothat Ti is not lost as TiO₂. Calcium treatment of oxide inclusions wasoptimized to modify alumina as well as sulfide inclusions. Low nitrogencontent Fe-alloys were used to result in low nitrogen in the melt. Lowernitrogen helps to control the amount of titanium lost to nitrogen andthus aids in enhancing the Ti:C stoichiometry. The heats were cast atcontinuous caster with suitable mold powder to result in good surfacequality of hot rolled coils. The steels were cast in to slabs of 55-65mm thickness and fed continuously to 6-stand hot strip processing millthrough a 290 meter long tunnel furnace maintained at temperature toeject slabs at exit temperatures of 1120-1130° C. Thin slabs (≤65 mm),fast casting speeds (5-5.5m/min) and low soaking temperature (1125° C.)helped achieving finer starting austenite grains than in a thick slabcasting unit.

The slabs were hot rolled to 2.2-3.3 mm thickness using a six-passreduction schedule and finish hot deformation temperature of 900-925° C.The hot rolled strip was immediately accelerated cooled to a coilingtemperature of 600-675° C. A cooling rate of more than 30° C./s wasemployed using super reinforced laminar cooling to result innon-polygonal ferrite grains with substructures if possible. Hot rolledsamples were collected from 600 cm inside of the outer lap of coilsafter cooling to evaluate mechanical properties and precipitation.

Cold Rolling and Choice of Annealing Parameters

The hot rolled coils were cold rolled to 0.85-1.2 mm thickness (60-65%cold reduction) and then annealed in a batch annealing furnace. Theannealing cycle to be used for these cold worked coils were initiallydetermined using laboratory simulations of batch annealing. Stripsamples prepared from cold rolled steels of both types were subjected toa programmed batch annealing cycle at various temperatures in a boxfurnace and tested after furnace cooling for mechanical properties andmicrostructure. FIG. 4 shows the variation in yield strengths as afunction of simulated batch annealing temperature. From the simulationresults it was evident that a batch annealing temperature window of600-650° C. was appropriate in achieving yield strength of 490 MPa andabove in both the steels. Batch annealing cycle was utilized to fullyrecover and completely recrystallize the cold worked ferrite grains.

FIG. 4 shows the yield and tensile properties of cold worked samples ofselected steels after simulated batch annealing. Based on thissimulation result cold rolled coils were then batch annealed in ahydrogen annealing furnace at temperatures between 600-650° C. Aftercooling, samples were collected from both steel grades from head andtail ends for mechanical and microstructural property evaluation.

Mechanical and Microstructural Properties Evaluation

Samples from batch annealed coils were collected from both head and tailends for various property evaluations such as hardness, tensileproperties and microstructures. Tensile samples across the width werealso tested to check cross width property variation.

Microstructural Characterization

Full thickness section along rolling direction was examined in opticalmicroscope (Leica DMI5000-M) as well as in scanning electron microscope(Hitachi SU3500) for microstructural features and cleanliness study.Metallographic samples were mechanically polished to 1 μm diamond pastesuspension followed by surface treatment in a Hitachi IM4000 ion millingsystem for orientation imaging mapping. An area of 128 mm×100 mm with astep size of 0.3 μm was chosen for EBSD analysis for orientation imagingand texture evaluation in a Hitachi SU3500 SEM.

Formability Evaluations

Hole expansion ratio tests were done on annealed sheet samples from bothsteels using an Interlaken SP400 test equipment at AMT-Fadi LLC. For thehole expansion tests, five square coupons of 100 mm×100 mm size were cutfrom each of three locations across the width-quarter width, center andthree-quarter width locations for examining property homogeneity. Holesof 10 mm diameter (d_(o)) was punched at center of each coupon for HERtests.

The test coupons were clamped between a holder and die with a clampingforce of 100 kN. A clearance of 12±1% of nominal sheet thickness wasadopted conforming to ISO 16630:2009(E) specification. A conical punchwith 60° angle was pierced through the hole at a speed of 0.25 mm/s andthe crack appearance during piercing was monitored using digital imagingsystem. The piercing was done at least after 30 mins after punching thehole. Diameter of the holes after crack appearance was measured and holeexpansion ratio, λ was calculated as

${\lambda = {\frac{\left( {d_{f} - d_{o}} \right)}{d_{o}} \times 100}},$

where d_(o) and d_(f) are initial and final diameter of the holerespectively.

Precipitate Examination in TEM

Precipitation in hot rolled samples as well as final fully processedsheet samples were studied in Talos L120C and Talos 200X electronmicroscope equipped with EDS X-ray spectrometer and electron-energy-lossspectroscopy (EELS). Samples of dimension 10 mm×10 mm were used toprepare TEM foils and replicas using a precision cutter from the centerareas of the sample parallel to rolling direction. For foilspreparation, samples were thinned by careful mechanical grinding andpolishing down to a thickness of around 80 mm. 3 mm diameter discs werepunched from thinned sheet followed electro-polishing in electrolyte of10% perchloric acid in methanol at −40C and 16V. Additionally,extraction replicas were prepared from the same samples for fineprecipitation analysis.

Results

Tensile properties of batch annealed sheet steels of both steel typesare shown in Table 3. Table 3 shows that both the sheet steels metminimum yield strength of 490 MPa in fully annealed condition. Excellenttotal elongation values were obtained. Ti-steel sheets showed slightlyhigher elongation values compared to Ti—Nb steel. Tensile propertieswere also outstandingly uniform from head to tail of the coils as wellas across the width of the coils as indicated in FIG. 5 . Table 3 showsthe tensile properties of batch annealed coils.

TABLE 3 Thickness, mm Yield Strength, Tensile Strength, SAE 490X MPa MPa% El Steel Location Spec 490-590 560 Min 14 Min n (10-e_(u)) Ti Head0.85-1.2 506-536 571-619 20-22 0.13-0.16 Tail 0.85-1.2 511-554 587-62820-24 0.13-0.14 Ti + Nb Head 0.85-1.2 512-546 596-616 19-23 0.14-0.15Tail 0.85-1.2 509-544 594-622 21-23 0.14-0.14

FIG. 5 shows the variation in strength properties of Ti- and Ti—Nbsteels across the coil width after batch annealing.

Microstructure

FIGS. 6 a and 6 b show the through thickness microstructures of hotrolled Ti-steel and Ti—Nb steel along the rolling direction. Ferritegrains in Ti-steel indicated significant non-polygonalilty due toaccelerated cooling from finishing temperature. Microstructure of Ti—Nbsteel represented finer and elongated ferrite grains. Ferrite grains areslightly coarser (d_(av)˜4.2 mm) in Ti-steel compared to Ti—Nb steel(d_(av)˜3.2 mm). The elongated ferrite grain structure in Ti—Nb steelwere indicative of partial non-recrystallization of austenite duringfinishing rolling whereas Ti-steel represented a fully recrystallizedmicrostructure prior to onset of accelerated cooling.

FIGS. 6 a and 6 b show the optical microstructures of Ti steel and Ti—Nbsteel after hot rolling. FIGS. 7 a through 7 d show the optical andscanning electron micrographs of batch annealed (a), (c) Ti-steel, and(b), (d) Ti—Nb steel showing fully recrystallized microstructures.

FIGS. 7 a through 7d revealed microstructures of fully processedTi-steel and Ti—Nb steel after batch annealing. Ferrite grains were muchrefined and nearly equiaxed in Ti-steel. Ti—Nb steel also represented afully recrystallized microstructure and ferrite grains are slightlyelongated. Scanning electron images of both steels indicated furtherclarity on the completion of recrystallization and onset of ferritegrain growth.

Cleanliness of the steels were evaluated through inclusion mapping inscanning electron microscope with EDS. Ternary diagram of oxide andsulfide inclusions were plotted and inclusion area fraction of variousprominent inclusion types were measured. FIG. 8 a shows typicalinclusions present in both steels revealing only globular oxides andoxy-sulfide inclusions. SEM-EDS microanalysis of globular inclusionsindicated these to be primarily CaO—Al₂O₃ and CaO—Al₂O₃—CaS inclusions.A typical SEM-EDS mapping of such inclusions is shown in FIG. 8revealing elemental presence of Ca, Al, O and S. Summary of significantinclusion types are shown in Table 4 indicating excellent cleanliness.

Table 4 shows a summary of area fraction of various inclusions presentin both steels as evaluated through SEM-EDS microanalysis.

TABLE 4 Area Fraction Inclusion Type Ti-Steel Ti—Nb-Steel Oxides(Ca-aluminate, Alumina, 1.56 × 10⁻⁵ 3.45 × 10⁻⁵ Spinel etc.) Sulfides(MnS, CaS, CaMnS 0 2.68 × 10⁻⁶ Oxy-sulfides (Ca-aluminate-CaS, 2.92 ×10⁻⁵  4.0 × 10⁻⁵ Ca-aluminate-CaMnS)

Precipitation Studies in Hot Rolled and Annealed Samples Ti—Nb Steel

In general, both coarse and fine nano-sized precipitates were seen inthe hot rolled and annealed specimens. Most of the coarse particles inhot rolled sample were 10-80 nm in size with an average size of 31.2±1.7nm and seen in low magnification in thin foil samples in FIG. 9(a).Coarse precipitates were seen mostly at grain boundaries as shown inFIG. 9 a . FIG. 9 b shows an overview of fine precipitates of less than8 nm in size and were distributed homogenously within ferrite. Fineprecipitates were also found in local areas as interphase precipitateswith row arrangements (FIG. 9 c ). The shape of the fine precipitateswas quasi-square and round.

FIG. 9 a shows TEM images of hot rolled Ti—Nb steel showing coarseprecipitates along grain boundaries. FIG. 9(b) is an overview of fineprecipitates and (c) some localized areas of precipitates in rowarrangement.

Chemical compositions of the ppts were analyzed using replica samplesusing EELS and EDS. EELS was performed by Digital Micrograph programfrom Gatan/Ametek and EDS was performed by Velox from Thermo-fisherScientific. A majority of precipitates were less than 20 nm were Ti andNb containing as shown in FIG. 10 a . No nitrogen was detected in them.These were (Ti,Nb)C precipitates and possibly formed during hot rolling.Particles coarser than 20 nm were mainly Ti4C2S2 or TiN inclusions. Highangle annular dark field image (HAADF) and EDS maps of particles lessthan 10 nm indicated Ti as major concentration (FIG. 10 b ).

FIGS. 10 a and 10 b show the chemical composition of precipitates in hotrolled Ti—Nb steel. Majority of ppts less than 20 nm (a) were Ti and Nbcontaining and no nitrogen was detected. (b) HAADF and EDS maps ofprecipitates less than 10 nm in size indicating Ti as major element.

Precipitate analysis in fully processed annealed sheet samples of Ti—Nbsteel revealed similar size, distribution of precipitates. Coarseprecipitates of 20-50 nm in sizes were Ti and Nb rich. Fine precipitatesof 3-11 nm were found distributed homogenously in ferrite. Ti was themain element detected in the fine particles as seen in FIG. 11(b).

FIGS. 11 a and 11 b show (a) an overview of fine nano-sized precipitateswithin ferrite grains and some along the grain boundaries. (b) HAADFimages and EDS mapping of precipitates 3-11 nm in size in fully annealedTi—Nb steel.

Precipitates in Ti-Steel

FIGS. 12 (a)-(b) show the precipitate distribution in hot rolled thinfoil sample. Precipitates were homogenously distributed within ferritegrains and some precipitation could be observed along grain boundaries.No row arrangement of precipitates or interphase precipitates were seenin Ti-steel. Coarse precipitates were fewer than Ti—Nb steel and sizedistribution was narrower between 20-30 nm.

These particles were mostly characterized as Ti, S bearing as shown inHAADF image and EDS mapping in FIG. 13 and possibly formed duringcasting. Fine nano-sized precipitates with size ranging from 1.5 nm to8.8 nm were found distributed uniformly within ferrite grains with anaverage of 4.3 nm as shown in FIG. 11(c). In order to identify the fineprecipitates extraction replica technique was used and identified byEELS. HAADF image and EDS mapping as shown in FIG. 14(a) and STEM imageand EELS maps identify them all fine precipitates as TiC precipitates(FIG. 14(b)).

FIG. 12 shows an overview of precipitates in hot rolled Ti-Steel asexamined in thin foils, (C) distribution of fine precipitates. FIG. 13shows an HAADF image and EDS mapping of precipitates larger than 20 nmwere Ti, C and S bearing.

FIG. 14 shows HAADF image and EDS mapping of nano-sized precipitates inhot rolled Ti-steel indicating Ti as major element. FIG. 15 shows EELSspectrum showing the presence of Ti and carbon only in a fineprecipitate.

Precipitate analysis in fully processed annealed sheet samples ofTi-steel revealed very fine nano-sized precipitates of 3-11 nm in sizedistributed homogenously in ferrite. Ti was the main element detected inthe fine particles as seen in FIG. 15(a). All precipitates below 15 nmwere identified as TiC (FIG. 15 b ).

FIGS. 16 a and 16 b show (a) Fine precipitates in ferrite in Ti-steel(b) HAADF image and EDS mapping of precipitates in Ti-steel revealingpresence of Ti only in fine precipitates less than 10 nm.

Formability Evaluation: Hole Expansion Ratio Results

Table 5 lists the hole expansion ratio (HER) values obtained for boththe Ti-steel and Ti—Nb-steel in fully processed condition. HER data fromvarious locations across the width were obtained and an average of fivesamples from each location are summarized. As indicated from theresults, outstanding HER values were obtained for both the steels. Thevalues are very uniform across the width of the coils. The HER valuesare indicative of excellent edge ductility or stretchability of thesteel so developed and is expected to perform well during stampingoperations. Table 5 shows the hole expansion ratio testing resultsperformed on annealed HSLA sheet steels.

TABLE 5 Thick- Yield HER Values λ, % ness, Strength, Quarter CenterQuarter Steel mm MPa width width Width Ti 1.2 501 96.7 ± 5.0  108.0 ±6.7 103.4 ± 7.4 Ti + 1.2 517 90.2 ± 12.7  91.9 ± 10.1  89.3 ± 7.2 Nb

Both Ti-only and Ti—Nb bearing HSLA steels of the selected lean alloycompositions successfully yielded a yield strength minimum of 490 MPa incold rolled batch annealed condition. Nb microalloying did influence theaustenite grain size after finishing deformation as could be seen fromelongated ferrite grain structure (FIG. 6 b ) and presence of grainboundary (Ti,Nb)C precipitates (FIG. 9 a ). In contrast, Ti-steelrepresented a fully recrystallized microstructure after finishdeformation rolling and was manifested in final ferrite grain structure.Non-recrystallization temperature, T_(NR) for Ti-steel was estimatedmuch lower (900° C.) than the finish deformation temperature employedcompared to Ti—Nb steel (967° C.) as estimated using experimentalevaluations. There was no advantage of adding Nb in the final coldrolled and annealed microstructure as the grain size advantage in hotrolled steel was later marginalized after cold rolling andrecrystallized annealing. Review of tensile properties in batch annealedsheet steels indicated Ti-steel indicated a marginal higher elongationvalues and better tensile properties across the width of the coils.

FIG. 17 shows the TEM images of fully recrystallized microstructures ofTi and (b) Ti—Nb steels. Hot rolling processing parameters wereoptimized for achieving maximum precipitation strengthening in the hotrolled coils. Batch annealing parameters after cold rolling wereoptimized to attain full recrystallization annealing before ferritegrain growth commences. Softening behavior of cold worked samples asstudied through batch annealing simulation (FIG. 4 ) indicated anoperable temperature window between 600-650° C. for the selected alloycomposition. Beyond this temperature ferrite grain coarsening observedand a substantial loss in strength detected as shown in ferritemicrostructures in FIG. 18 . FIG. 18 shows the microstructures ofTi-steel after simulated batch annealing at (a) 630° C. and (b) 700° C.revealing significant ferrite grain growth after recrystallization.

Both steels represented similar softening behavior with temperature.Optical micrographs as well as transmission electron micrographs ofannealed samples did indicate full recrystallization of ferrite grains.(FIG. 6 , FIG. 17 ). No substructures were observed in either of thesteels. Ti-steel manifested a slightly finer final grain structure thanthe Nb-added Ti-steel. Further studies of texture evaluation and grainboundary orientation of ferrite grains were carried out using electronback scattered diffraction (EBSD).

FIG. 19 shows the grain boundary mapping of (a) Ti-steel and (b) Ti—Nbsteel constructed from orientation imaging indicating significantfraction of high angle grain boundary (HAGB) (˜79-83%) in both steelsrevealing near complete recrystallization in both steels. Blue linesrepresent high angle grain boundaries with critical misorientation of15o or red lines representing low angle boundaries with misorientationof 2 to 15o. Grain size distributions as calculated using grain boundarymapping is shown in (c).

Grain size distribution as estimated from grain boundary mapping isplotted in FIG. 18(c) and an average ferrite grain size of 3.9 mm and4.6 mm were estimated for Ti-steel and Ti—Nb steel respectively.

The annealing textures of both Ti-steel and Ti—Nb steel are representedin FIGS. 20 a and 20 b and showed similar recrystallization textures.Both steels revealed increasing γ-fiber with presence of some α-fiber aswell. The γ-fiber showed significant intensity concentration or maximumat {111}<112> positions albeit non-uniform along the γ-fiber location.Ti-steel revealed relatively a stronger γ-fiber alignment compared toTi—Nb steel. The ferrite grain boundary orientation indicated in FIG. 19revealed a mixed grain structure which is probably the reason for notobtaining a sharp and uniform γ-fiber texture. Sharper γ-fiber textureis typical of extra deep drawing quality steels with a coarser andequiaxed ferrite grain structure. FIG. 20 shows the ODF of batchannealed Ti-steel and Ti—Nb steel showing increasing g-fiberorientations.

Precipitation studies revealed significant nano-sized TiC precipitates(1.5 nm to 8 nm) in both steels strongly contributing towards the yieldstrength increment. The precipitates that formed after hot rollingremained mostly intact in final batch annealed steels as the size anddistribution of fine precipitate did not change during recrystallizationannealing. In Ti-steel most precipitates were finer than 10 nm in sizeand lesser precipitates coarser than 10 nm were found. In contrast,Ti—Nb steel showed relatively coarser (Ti,Nb)C precipitates. FIG. 21shows the precipitate evolution in hot rolled to cold rolled andannealed microstructure. Distribution and average size of precipitateswere mostly similar. Average TiC precipitate size could be evaluated as3.16 nm and 3.96 nm in Ti and Ti—Nb steels respectively. These nanoprecipitates could successfully contribute to the yield strength of theferrite. Noting an effective Ti available for precipitation in Ti-steeland Ti—Nb steel as 0.067 wt. % and 0.058 wt. % (after nitrogenstabilization and Nb precipitation) the strengthening incrementcalculated using Eq. [1] can be 102 MPa and 87 MPa respectively.

Hole expansion ratio data obtained in the current batch annealed steelindicated superior stretchability or edge ductility suitable for moststamping operations. In general, both sheet steels indicated high HERvalues because of internal cleanliness, finer ferrite grain size andnano precipitates. Higher HER values obtained for Ti-alloyed HSLA steelwas possibly due to absence of grain boundary coarser (Ti,Nb)Cprecipitates and relatively finer annealed ferrite grain structure thanNb added steel.

In some embodiments, hole expansion ratio values and tensile propertiesof the steel provides steel that can be utilized for autobody stampingapplications requiring high stretch ductility.

Notwithstanding that the numerical ranges and parameters setting forththe broad scope of the invention are approximations, the numericalvalues set forth in the specific examples are reported as precisely aspossible. Any numerical value, however, inherently contains certainerrors necessarily resulting from the standard variation found in theirrespective testing measurements.

Any numerical range recited herein is intended to include all sub-rangessubsumed therein. For example, a range of “1 to 10” is intended toinclude all sub-ranges between (and including) the recited minimum valueof 1 and the recited maximum value of 10, that is, having a minimumvalue equal to or greater than 1 and a maximum value of equal to or lessthan 10.

In this application, the use of the singular includes the plural andplural encompasses singular, unless specifically stated otherwise. Inaddition, in this application, the use of “or” means “and/or” unlessspecifically stated otherwise, even though “and/or” may be explicitlyused in certain instances. In this application and the appended claims,the articles “a,” “an,” and “the” include plural referents unlessexpressly and unequivocally limited to one referent.

As used herein, “including,” “containing” and like terms are understoodin the context of this application to be synonymous with “comprising”and are therefore open-ended and do not exclude the presence ofadditional undescribed or unrecited elements, materials, phases, ormethod steps. As used herein, “consisting of” is understood in thecontext of this application to exclude the presence of any unspecifiedelement, material, phase, or method step. As used herein, “consistingessentially of” is understood in the context of this application toinclude the specified elements, materials, phases, or method steps,where applicable, and to also include any unspecified elements,materials, phases, or method steps that do not materially affect thebasic or novel characteristics of the disclosure.

1. A high strength low alloy steel sheet product produced by acontinuous strip process, the sheet product comprising from 0.045 to0.06 weight percent Carbon, from 0.75 to 1.2 weight percent Manganese,from 0.02 to 0.04 weight percent Aluminum, and from 0.075 to 0.12%weight percent Titanium and an Fe balance, wherein the steel sheetproduct has been subjected to cold rolling and batch annealing to form asteel sheet product having a yield strength of at least 500 MPa, asubstantially ferritic microstructure with nano TiC precipitates and ahole expansion ratio of at least 60%.
 2. The high strength low alloysteel sheet product of claim 1, wherein the TiC precipitates areapproximately 3.0 to 7.0 nm.
 3. The high strength low alloy steel sheetproduct of claim 1, wherein the steel sheet comprises 0.06 weightpercent Carbon, 1.1 weight percent Manganese and 0.12 weight percentTitanium.
 4. The high strength low alloy steel sheet product of claim 1,wherein the steel sheet comprises a hole expansion ratio of at least70%.
 5. The high strength low alloy steel sheet product of claim 1further comprising a tensile strength of at least 570 MPa.
 6. The highstrength low alloy steel sheet product of claim 1 further comprisingless than 0.04 weight percent Niobium.
 7. The high strength low alloysteel sheet product of claim 1 wherein the steel sheet is produced froma slab of approximately 55 mm to 85 mm thick.
 8. The high strength lowalloy steel sheet product of claim 1, wherein the steel sheet isproduced with a coiling temperature of 600 to 675C.
 9. The high strengthlow alloy steel sheet product of claim 1, wherein the steel sheet isproduced by cold rolling to 40 to 75%.
 10. The high strength low alloysteel sheet product of claim 1, wherein the steel sheet is produced hotrolling the slabs to a 2.2-3.3 mm thickness using a six-pass reductionschedule.
 11. The high strength low alloy steel sheet product of claim 1wherein the steel sheet is produced by hot rolling the steel slab at ahot rolling finishing temperature of 875 to 950C.
 12. The high strengthlow alloy steel sheet product of claim 11, wherein the hot rolled stripis cooled at a cooling rate of at least 30° C./s to a coilingtemperature of 600-675° C.
 13. A method for producing a high strengthlow alloy steel sheet product comprising from 0.045 to 0.06 weightpercent Carbon, from 0.75 to 1.2 weight percent Manganese, from 0.02 to0.04 weight percent Aluminum, and from 0.075 to 0.12% weight percentTitanium and an Fe balance, a yield strength of at least 500 MPa, asubstantially ferritic microstructure with nano TiC precipitates and ahole expansion ratio of at least 60%, comprising the steps of:continuously casting a steel slab approximately 55 mm-85 mm thick;maintaining a temperature of the steel slab; hot rolling the steel slabto a steel sheet at a finishing temperature of 875-950C; cooling thesteel slab; coiling the steel sheet at a temperature of 600-675C; coldrolling the steel sheet to 40-75%; and hydrogen batch annealing thesteel sheet at temperature of 600-650C to achieve a fully recrystallizedferritic microstructure with nano precipitates of TiC.
 14. The methodfor producing a high strength low alloy steel sheet product of claim 13,wherein the step of hot rolling includes hot rolling the slabs to a2.2-3.3 mm thickness using a six-pass reduction schedule.
 15. The methodfor producing a high strength low alloy steel sheet product of claim 13,wherein the step of cooing the steel slab includes a cooling rate ofmore than 30° C./s.
 16. The method for producing a high strength lowalloy steel sheet product of claim 13, wherein the steel sheet includesa tensile strength of at least 570 MPa.
 17. The method for producing ahigh strength low alloy steel sheet product of claim 13, wherein thesteel sheet includes a hole expansion ratio of least 70%.
 18. A highstrength low alloy steel sheet product produced by a continuous stripprocess, the sheet product comprising from 0.045 to 0.06 weight percentCarbon, from 0.75 to 1.2 weight percent Manganese, from 0.02 to 0.04weight percent Aluminum, from 0.075 to weight percent Titanium, 0.04weight percent Niobium and an Fe balance, wherein the steel sheetproduct has been subjected to cold rolling and batch annealing to form asteel sheet product having a yield strength of at least 500 MPa, asubstantially ferritic microstructure with nano TiC precipitates and ahole expansion ratio of at least 60%.
 19. The high strength low alloysteel sheet product of claim 18, wherein the TiC precipitates areapproximately 3.0 to 7.0 nm.
 20. The high strength low alloy steel sheetproduct of claim 18, wherein the steel sheet comprises 0.06 weightpercent Carbon, 1.1 weight percent Manganese and 0.12 weight percentTitanium.
 21. A high strength low alloy steel sheet product produced bya continuous strip process, the sheet product comprising from 0.045 to0.06 weight percent Carbon, from 0.75 to 1.2 weight percent Manganese,from 0.02 to 0.04 weight percent Aluminum, and from 0.075 to 0.12%weight percent Titanium and an Fe balance, wherein the steel sheetproduct has been subjected to cold rolling and batch annealing to form asteel sheet product having a yield strength of at least 550 MPa, asubstantially ferritic microstructure with nano TiC precipitates and ahole expansion ratio of at least 60%.
 22. The high strength low alloysteel sheet product of claim 21, wherein the TiC precipitates areapproximately 3.0 to 7.0 nm.
 23. The high strength low alloy steel sheetproduct of claim 21, wherein the steel sheet comprises 0.06 weightpercent Carbon, 1.1 weight percent Manganese and 0.12 weight percentTitanium.
 24. The high strength low alloy steel sheet product of claim21, wherein the steel sheet comprises a hole expansion ratio of at least70%.